Method of improving the mechanical properties of a component

ABSTRACT

A method ( 40 ) of improving the mechanical properties of a component, for example a gas turbine engine turbine disc, ( 24 ) comprises isothermally forging ( 42 ) a preform to produce a shaped preform with a predetermined shape at a first predetermined temperature, solution heat treating ( 44 ) the shaped preform, quenching ( 46 ) the shaped preform, forging ( 48 ) the shaped preform at a second predetermined temperature to impart a predetermined residual strain in the shaped preform, ageing ( 50 ) the shaped preform and finally machining ( 52 ) the shaped preform to a finished shape. The second predetermined temperature is less than the first predetermined temperature.

The present invention relates to a method of improving the mechanicalproperties of a component, in particular to a method of improving themechanical properties of a forged nickel base superalloy article, e.g. aforged nickel base superalloy gas turbine engine turbine disc or aforged nickel base superalloy gas turbine engine compressor disc.

High strength nickel base superalloys for critical rotor components,e.g. turbine discs or compressor discs, are made by complex powderprocessing route. Other high strength nickel base superalloys forcritical rotor components are made by a cast and wrought route.

The complex powder processing route comprises vacuum induction melting(VIM) of the nickel base superalloy, inert gas, e.g. argon, atomisation(IGA) to produce a metal powder, sieving of the metal powder, blendingof the metal powder, canning of the metal powder, hot isostatic pressing(HIP) of the metal powder, extrusion to form a billet, isothermalforging of the billet to form a forging, solution heat treatment (SHT)of the forging, ageing heat treatment (AHT) of the forging and machiningthe forging to form the final shape of a component.

The cast and wrought route comprises vacuum induction melting (VIM),electro slag refining (ESR), vacuum arc remelting (VAR), conversion ofthe ingot to a billet through multiple upset and heat treatmentoperations, isothermal forging of the billet to a forging, solution heattreatment (SHT) of the forging, ageing heat treatment (AHT) of theforging and machining the forging to form the final shape of acomponent.

Lower strength and/or lower temperature capability nickel basesuperalloys, such as Waspaloy or IN718, forgings are formed to shapeusing conventional forging, e.g. press forging. The forgings undergo acomplex heat treatment cycle to optimise the grain size and thestrengthening precipitate phase distribution. The forging is thenmachined to the final shape of a component.

Currently higher strength and/or higher temperature capability nickelbase superalloys are optimised to provide maximum creep and fatiguecrack growth properties and/or maximum tensile and fatigue properties byaltering the grain size of the nickel base superalloy and the volumefraction and size of strengthening precipitates, which results in atrade off in mechanical properties. Therefore, it is desirable toincrease the creep properties and/or tensile properties of a highstrength nickel base superalloy without altering the grain size of thenickel base superalloy or the size and/or distribution of otherstrengthening phases in the nickel base superalloy.

Residual stresses develop in components due to the thermal gradientsthat develop during the cooling steps of a heat treatment cycle.Solution heat treatments which define the grain size in these componentsare followed by quenching in air, oil or other medium. The quenching isto minimise the size of the precipitate phase that is responsible forhigh temperature strength in these nickel base superalloys. Byoptimising the precipitate size the component is subjected to a largethermal gradient during quenching and this thermal gradient produceslarge residual stresses in the component. The residual stresses may leadto distortion of the component during subsequent machining. In practicethe cooling rates during quenching are reduced to avoid quench cracking.Nickel base superalloys have an ageing heat treatment after the solutionheat treatment to optimise the precipitates further and to relieveresidual stresses in the component.

In nickel base superalloys used as turbine discs, or compressor discs,the residual stresses in the disc are added to the stresses induced inthe disc by the operation of the gas turbine engine. This combinedstress must be maintained below the stress level that would be predictedto cause tensile or fatigue failure of the turbine disc or compressordisc. Thus, the residual stress in the turbine disc, or compressor disc,limits the mechanical stress cycle that the engine is designed for.

Accordingly the present invention seeks to provide a novel method ofmanufacturing a component which reduces, preferably overcomes, the abovementioned problem or problems.

Accordingly the present invention provides a method of improving themechanical properties of a component comprising the steps of: —a)forging a preform to produce a shaped preform with a predetermined shapeat a first predetermined temperature,

b) solution heat treating the shaped preform,c) quenching the shaped preform,d) ageing the shaped preform,e) machining the shaped preform to a finished shape or a semi-finishedshape, andf) forging the shaped preform at a second predetermined temperature toimpart a predetermined residual strain in the shaped preform after stepc) and before step e), wherein the second predetermined temperature isless than the first predetermined temperature.

Step f) may be after step c) and before step d), step f) may be afterstep d) and before step e) or step f) may be concurrent with step d).

Preferably the second predetermined temperature is between 700° C. and870° C.

More preferably the second predetermined temperature is between 750° C.and 850° C. More preferably the second predetermined temperature isbetween is 760° C. and 810° C. The second predetermined temperature maybe 760° C., 802° C. or 843° C.

Preferably the forging step f) imparts a predetermined residual tensilestrain or a predetermined residual compressive strain.

Preferably the forging step f) imparts a strain of less than 10%.

Preferably step f) comprises isothermally forging.

Preferably step f) comprises forging at a strain rate between 1×10⁻⁴ and1×10⁻² s⁻¹.

Preferably step a) comprises isothermally forging.

In step a) the first predetermined temperature may be up to gamma primesolvus minus 25° C. to 50° C. In step a) the forging may be at a strainrate between 1×10⁻⁴ and 1×10⁻² s⁻¹.

Alternatively in step a) the first predetermined temperature may be upto gamma prime solvus minus 55° C. to 110° C. In step a) the forging maybe at a strain rate between 1×10⁻² and 5×10⁻¹ s⁻¹.

Preferably the method comprises machining the shaped preform after stepa) and before step b).

Step b) may comprise a subsolvus solution heat treatment and or asupersolvus heat treatment,

Step b) may comprise a subsolvus solution heat treatment at 1120° C. for4 hours.

Step d) may comprise an ageing heat treatment at 760° C. for 16 hours.

Step b) may comprise a subsolvus solution heat treatment at 1120° C. for4 hours, followed by quenching, followed by a supersolvus heat treatmentat 1204° C. for 1 hour.

Step b) may comprise a supersolvus heat treatment at 1204° C. for 1hour.

Preferably the component is a compressor disc, a turbine disc, acompressor cone or a turbine cover plate.

Preferably the component comprises a nickel base superalloy or atitanium base alloy.

The nickel base superalloy may be RR1000, U720Li, Rene 95, Rene 88DT,ME3, N18, Alloy 10, LSHR and other nickel base superalloys suitable forapplication as a turbine disc or compressor disc.

The preform may have been made by a cast and a wrought route oralternatively may have been made by a powder processing route.

The present invention will be more fully described by way of examplewith reference to the accompanying drawings in which:—

FIG. 1 shows a gas turbine engine having a turbine disc which has beenmanufactured according to the present invention.

FIG. 2 shows is a cross-sectional view through a portion of a gasturbine engine turbine disc which has been manufactured according to thepresent invention.

FIG. 3 is a flow chart of a method of manufacturing a componentaccording to the present invention.

FIG. 4 is a flow chart of a further method of manufacturing a componentaccording to the present invention.

FIG. 5 is a flow chart of another method of manufacturing a componentaccording to the present invention.

FIG. 6 is a bar chart showing the ultimate tensile strength and the 0.2%proof strength in tensile tests for fine grained RR1000 processedconventionally and according to the present invention.

FIG. 7 is a bar chart showing the percentage elongation and thepercentage reduction in area in tensile tests for fine grained RR1000processed conventionally and according to the present invention.

FIG. 8 is a bar chart showing the ultimate tensile strength and the 0.2%proof strength in tensile tests for coarse grained RR1000 processedconventionally and according to the present invention.

FIG. 9 is a bar chart showing the percentage elongation and thepercentage reduction in area in tensile tests for coarse grained RR1000processed conventionally and according to the present invention.

FIGS. 10A, 10B and 10C are graphs showing the hoop stress, radial stressand axial stress in test pieces at different axial and radial positionsin a first cylindrical test piece which was water quenched only.

FIGS. 11A, 11B and 11C are graphs showing the hoop stress, radial stressand axial stress in test pieces at different axial and radial positionsin a second cylindrical test piece which was water quenched, aged andgiven a low deformation.

FIGS. 12A, 12B and 12C are graphs showing the hoop stress, radial stressand axial stress in test pieces at different axial and radial positionsin a third cylindrical test piece which was water quenched and given ahigh deformation.

FIGS. 13A, 13B and 13C are graphs showing the hoop stress, radial stressand axial stress in test pieces at different axial and radial positionsin a fourth cylindrical test piece which was water quenched and given alow deformation.

FIGS. 14A, 14B and 14C are graphs showing the hoop stress, radial stressand axial stress in test pieces at different axial and radial positionsin a fifth cylindrical test piece which was oil quenched, aged and givena medium deformation.

FIGS. 15A, 15B and 15C are graphs showing the hoop stress, radial stressand axial stress in test pieces at different axial and radial positionsin a sixth cylindrical test piece which was oil quenched and aged.

A turbofan gas turbine engine 10, as shown in FIG. 1, comprises in axialflow series an intake 12, a fan section 14, a compressor section 16, acombustion section 18, a turbine section 20 and an exhaust 22. Theturbine section 20 comprises a turbine disc 24, which carries aplurality of circumferentially space turbine blades 26. The gas turbineengine is quite conventional and its construction and operation will notbe described further.

The gas turbine engine turbine disc 24, as shown more clearly in FIG. 2,comprises a hub, or cob, 26, a web 28 and a rim 30. The hub 26 is at theradially inner end of the turbine disc 24, the rim 30 is at the radiallyouter end of the turbine disc 24 and the web 28 extends radially betweenand interconnects the hub 26 and the rim 30. The rim 30, in thisexample, has a plurality of circumferentially spaced slots 34 to receivethe roots of turbine blades 26, shown in FIG. 1, and circumferentiallyspaced posts 32 are provided on the rim 30 of the turbine disc 24 todefine the sides of the slots 34. The slots 34 may be firtree shape, ordovetail shape. The turbine disc 24 comprises a high strength nickelbase superalloy, for example RR1000.

A first method 40 of improving the mechanical properties of a component,for example a gas turbine engine turbine disc, 24 according to thepresent invention, as shown in FIG. 3, comprises isothermally forging 42a preform to produce a shaped preform with a predetermined shape at afirst predetermined temperature, solution heat treating 44 the shapedpreform, quenching 46 the shaped preform, forging 48 the shaped preformat a second predetermined temperature to impart a predetermined residualstrain in the shaped preform, ageing 50 the shaped preform and finallymachining 52 the shaped preform to a finished shape. It is to be notedthat the second predetermined temperature is less than the firstpredetermined temperature.

A second method 40B of improving the mechanical properties of acomponent, for example a gas turbine engine turbine disc, 24 accordingto the present invention, as shown in FIG. 4, comprises isothermallyforging 42 a preform to produce a shaped preform with a predeterminedshape at a first predetermined temperature, solution heat treating 44the shaped preform, quenching 46 the shaped preform, ageing 50 theshaped preform, forging 48 the shaped preform at a second predeterminedtemperature to impart a predetermined residual strain in the shapedpreform and finally machining 52 the shaped preform to a finished shape.It is to be noted that the second predetermined temperature is less thanthe first predetermined temperature.

A third method 40C of improving the mechanical properties of acomponent, for example a gas turbine engine turbine disc, 24 accordingto the present invention, as shown in FIG. 5, comprises isothermallyforging 42 a preform to produce a shaped preform with a predeterminedshape at a first predetermined temperature, solution heat treating 44the shaped preform, quenching 46 the shaped preform, simultaneouslyforging 48 the shaped preform at a second predetermined temperature toimpart a predetermined residual strain in the shaped preform and ageing50 the shaped preform and finally machining 52 the shaped preform to afinished shape. It is to be noted that the second predeterminedtemperature is less than the first predetermined temperature.

In the three methods discussed above the second predeterminedtemperature is between 700° C. and 870° C. (1300° F. and 1600° F.), morepreferably the second predetermined temperature is between 750° C. and850° C. (1380° F. and 1560° F.), even more preferably 760° C. to 810° C.(1400° F. to 1490° F.). The forging 48 step may be arranged to impart apredetermined residual tensile strain or a predetermined residualcompressive strain. The forging 48 step imparts a strain of less than orequal to 15%, e.g. 5% or 10% strain.

The three methods mention above may comprise machining the shapedpreform after the isothermal forging 42 and before the solution heattreatment 44.

Although the three methods mentioned previously have mentioned a gasturbine engine turbine disc, the component may be a compressor disc, acompressor cone or a turbine cover plate. The component may comprise anickel base superalloy, a titanium base alloy or other suitable alloy.

The preform used in the previously mentioned methods may have been madeby a cast and a wrought route or alternatively may have been made by apowder processing route.

The forging 48 may comprise isothermal forging, hot die press forging orhammer forging and the forging 48 may comprise applying a mechanicalload, a fluid load or a thermal gradient via any conventional forgingapparatus or process. The isothermal forging may use an isothermalforging press and the die for the isothermal forging press may compriseTZM molybdenum or other suitable material.

The final machining 52 may comprise any suitable machining, e.g.turning, grinding, milling, drilling, polishing etc.

The isothermal forging 42 is preferred, but may be replaced by othersuitable types of forging.

The present invention is described more fully with reference to anexample. RR1000 consists of 18.5 wt % cobalt, 15 wt % chromium, 5 wt %molybdenum, 2 wt % tantalum, 3.6 wt % titanium, 3 wt % aluminium, 0.5 wt% hafnium, 0.015 wt % boron, 0.06 wt % zirconium, 0.027 wt % carbon andthe balance nickel plus incidental impurities. RR1000 has a gamma primesolvus temperature of 1145° C. to 1150° C. Thus, a turbine, orcompressor, disc consisting of RR1000 is produced by initially producinga billet, using either powder metallurgy, or cast and wrought,techniques.

The RR1000 billet is then isothermally forged, at step 42 in FIG. 3, toproduce a shaped preform which is near to the final shape of the disc ata temperature up to gamma prime solvus minus 25° C. to 50°, at a strainrate between 1×10⁻⁴ and 1×10⁻² s⁻¹ or at a temperature up to gamma primesolvus minus 55° C. to 110° C. at a strain rate between 1×10⁻² and5×10⁻¹ s⁻¹. The RR1000 shaped preform is then solution heat treated, atstep 44, at a temperature in the range of gamma prime solvus minus 15°C. to 35° C. up to gamma prime solvus plus 25° C. to 60° C. for timesbetween 0.5 and 8 hours. The shaped preform is cooled or quenched, atstep 46, from the solution heat treatment temperature at a rate suitableto avoid quench cracking at stress concentrations, for example at a ratebetween 0.1° C. s⁻¹ and 10° C. s⁻¹. The shaped preform is thenisothermally forged, at step 48, at a temperature between 700° C. (1292°F.) and 870° C. (1598° F.), at a strain rate between 1×10⁻⁴ and 1×10⁻²s⁻¹ to impart a predetermined residual strain to the shaped preform. Theshaped preform is then given an ageing heat treatment, at step 50, at atemperature between 650° C. (1202° F.) and 800° C. (1472° F.) forbetween 2 and 30 hours. Finally the shaped preform is machined to finalshape at step 52.

A series of tests were carried out on samples of fine grained and coarsegrained RR1000 nickel base superalloy, which were initially forged.Samples 1 and 2 of RR1000 were given a conventional subsolvus solutionheat treatment at 1120° C. (2048° F.) for 4 hours, then air cooled,followed by an ageing heat treatment at 760° C. (1400° F.) for 16 hoursand then air cooled as a baseline. Other samples, samples 3 and 6, ofRR1000 were given a subsolvus solution heat treatment at 1120° C. (2048°F.) for 4 hours, then air cooled, followed by an ageing heat treatmentat 760° C. (1400° F.) for 16 hours and then strained at 760° C. (1400°F.) at 5% or 10% strain respectively. Other samples, samples 4 and 7, ofRR1000 were given a subsolvus solution heat treatment at 1120° C. (2048°F.) for 4 hours, then air cooled, followed by an ageing heat treatmentat 760° C. (1400° F.) for 16 hours and then strained at 802° C. (1475°F.) at 5% or 10% strain respectively. Other samples, samples 5, 22 and8, of RR1000 were given a subsolvus solution heat treatment at 1120° C.(2048° F.) for 4 hours, then air to cooled, followed by an ageing heattreatment at 760° C. (1400° F.) for 16 hours and then strained at 843°C. (1550° F.) at 5%, 10% or 15% strain respectively. Additional samples,samples 17 and 11, of RR1000 were given a subsolvus solution heattreatment at 1120° C. (2048° F.) for 4 hours, then air cooled, thenstrained at 760° C. (1400° F.) at 5% or 10% strain respectively,followed by an ageing heat treatment at 760° C. (1400° F.) for 16 hours.Another sample, sample 12, of RR1000 was given a subsolvus solution heattreatment at 1120° C. (2048° F.) for 4 hours, then air cooled, thenstrained at 802° C. (1475° F.) at 5% strain, followed by an ageing heattreatment at 760° C. (1400° F.) for 16 hours. Another sample, sample 13,of RR1000 was given a subsolvus solution heat treatment at 1120° C.(2048° F.) for 4 hours, then air cooled, then strained at 843° C. (1550°F.) at 5% strain, followed by an ageing heat treatment at 760° C. (1400°F.) for 16 hours.

Samples, samples 9 and 10, of RR1000 were given a conventional subsolvussolution heat treatment at 1120° C. (2048° F.) for 4 hours, then aircooled, followed by a supersolvus heat treatment at 1204° C. (2200° F.)for 1 hour, then air cooled, followed by an ageing heat treatment at760° C. (1400° F.) for 16 hours and air cooled as a baseline. Furthersamples, samples 14 and 19, of RR1000 were given a subsolvus solutionheat treatment at 1120° C. (2048° F.) for 4 hours, then air cooled,followed by a supersolvus heat treatment at 1204° C. (2200° F.) for 1hour, then air cooled, followed by an ageing heat treatment at 760° C.(1400° F.) for 16 hours and then strained at 760° C. (1400° F.) at 5% or10% strain respectively. Other samples, samples 15 and 20, of RR1000were given a subsolvus solution heat treatment at 1120° C. (2048° F.)for 4 hours, then air cooled, followed by a supersolvus heat treatmentat 1204° C. (2200° F.) for 1 hour, then air cooled, followed by anageing heat treatment at 760° C. (1400° F.) for 16 hours and thenstrained at 802° C. (1475° F.) at 5% or 10% strain respectively. Othersamples, samples 16 and 24, of RR1000 were given a subsolvus solutionheat treatment at 1120° C. (2048° F.) for 4 hours, then air cooled,followed by a supersolvus heat treatment at 1204° C. (2200° F.) for 1hour, then air cooled, followed by an ageing heat treatment at 760° C.(1400° F.) for 16 hours and then strained at 843° C. (1550° F.) at 5% or10% strain respectively. These samples were air cooled after thesupersolvus heat treatment at a rate of 0.81° Cs⁻¹. In all the abovesamples the samples were air cooled after the subsolvus heat treatmentat a rate of 0.76° Cs⁻¹.

In all cases the samples were held at the appropriate temperature for 1hour before any strain was applied.

The subsolvus heat treatment, followed by ageing heat treatment producedfine grains in the nickel base superalloy and the subsolvus heattreatment, followed by the supersolvus heat treatment and ageing heattreatment produced coarse grains in the nickel base superalloy as iswell known to those skilled in the art.

Then standard test pieces were taken from each of the large samples ofRR1000 and the test pieces of the samples were then subjected to tensiletests at a temperature of 650° C. (1202° F.) to determine the ultimatetensile strength and the 0.2% proof strength of the samples and todetermine the percentage elongation and percentage reduction in area ofthe samples. The results are recorded in Table A below and some of theresults are shown in FIGS. 6, 7, 8 and 9.

TABLE A Ultimate 0.2% Strain Tensile Proof Thermal Temp Strain StrengthStrength % % Red Sample History (° F.) (%) (MPa) (MPa) Elong Area 1 Sb +A — — 1382 1004 25 35 2 Sb + A — — 1381 1000 21 25 3 Sb + A + St 1400 51502 1211 12 37 4 Sb + A + St 1475 5 1484 1251 17 26 5 Sb + A + St 15505 1464 1176 22 36 6 Sb + A + St 1400 10 1598 1365 9 34 7 Sb + A + St1475 10 1503 1249 13 33 22 Sb + A + St 1550 10 1499 1265 16 38 8 Sb +A + St 1400 15.5 1487 1138 13 31 17 Sb + St + A 1400 5 1549 1277 15 2712 Sb + St + A 1475 5 1489 1238 15 34 13 Sb + St + A 1550 5 1462 1225 2133 11 Sb + St + A 1400 10 1498 1229 7 15 9 Sb + Su + A + St — — 1356 84622 25 10 Sb + Su + A + St — — 1365 852 22 24 14 Sb + Su + A + St 1400 51494 1209 9 18 23 Sb + Su + A + St 1550 5 1431 1129 13 27 19 Sb + Su +A + St 1400 10 1485 1240 4 12 20 Sb + Su + A + St 1475 10 1476 1218 9 1724 Sb + Su + A + St 1550 10 1541 1253 11 29 (Sb—subsolvus heattreatment, Su—supersolvus heat treatment, A—ageing heat treatment,St—strain heat treatment)

The above results show that the present invention has increased theultimate tensile strength and the 0.2% proof strength of a fine grainednickel base superalloy above that of a fine grained nickel basesuperalloy given a conventional subsolvus heat treatment followed by anageing heat treatment. The above results show that the present inventionhas increased the ultimate tensile strength and the 0.2% proof strengthof a coarse grained nickel base superalloy above that of a nickel basesuperalloy given a conventional subsolvus heat treatment, followed by asupersolvus heat treatment followed by an ageing heat treatment.

Another series of tests were carried out on samples of RR1000 nickelbase superalloy, which were initially forged. This series of tests usedthe method described with respect to FIG. 4. The quenching, step 46, ofthe nickel base superalloy was chosen to be either an oil quench or awater quench to impart a high level of strain into the isothermallyforged, step 42, and solution heat treated, step 44, nickel basesuperalloy. The ageing heat treatment, step 50, was a conventional ageat 760° C. (1400° F.) for 16 hours. The isothermal forging, step 48, wasconducted within the preferred temperature range at a temperature lessthan 760° C. (1400° F.).

In order to investigate the effect of strain, three different strainvalues were investigated. Residual stress was measured using a neutrondiffraction technique, which allows for non-destructive evaluation ofthe nickel base superalloys. The residual stresses were measured at anumber of locations in three different orientations, the orientationswere hoop, axial and radial. The test pieces were cylindrical andnominally had a diameter of 75 mm and a height, or thickness, of 25 mm.The residual stress was measured at locations at 5 mm, 12 mm and 19 mmheight from one face of the cylindrical test piece and at radiallocations of 0 mm, 10 mm, 19 mm, 30 and 33 mm from the centre of thecylindrical test piece. Table B shows the residual hoop stress levels,in MPa, at different radial locations at a height of 12 mm from thesurface of the cylindrical test pieces, for different quenching, ageingand deformation conditions. All the deformations are less than or equalto 10% strain.

TABLE B Test Deformation Height Radial Position (mm) Piece Quench Age(mm) 0 10 19 30 33 1 Water — — 12 1300 1400 1550 500 −300 2 Water AgedLow 12 675 750 850 350 0 3 Water — High 12 250 260 400 510 480 4 Water —Low 12 1000 1000 1200 600 −200 5 Oii Aged Medium 12 700 725 760 175 −2006 Oil Aged — 12 750 — 800 — −250Test piece 1, which was water quenched, but was not aged and was notdeformed is considered a baseline and it is seen that test piece 1 hashigh levels of residual stress present at all radial locations. Testpiece 2, which was water quenched, aged and given a low deformation hasa much lower levels of residual stress at all radial positions comparedto test piece 1 due to the combination of a conventional age and a lowdeformation and low temperature mechanical stress relief. Test piece 4,which was water quenched and given a low deformation has levels ofresidual stress intermediate that of test pieces 1 and 2 except for the30 mm radial position. Test piece 3, which was water quenched and givena high deformation has lower levels of residual stress than test piece 2at the 0 mm, 10 mm and 19 mm radial locations. Comparing test pieces 1,2 and 4 it can be seen that the deformation alone and the ageing anddeformation produce a reduction in the residual stress and thereforethat the combination of deformation and ageing produces a greaterreduction in the residual stress. Test piece 5, which was oil quenched,aged and given a medium deformation has lower levels of residual stressthan test piece 6, which was oil quenched and aged. FIGS. 10 to 15 aregraphs showing the hoop stress, radial stress and axial stress for testpieces 1, 2, 3, 4, 5 and 6 respectively at locations at 5 mm, 12 mm and19 mm axially, height, from one face of the cylindrical test piece andat radial locations of 0 mm, 10 mm, 19 mm, 30 and 33 mm from the centreof the cylindrical test piece. This data shows the effectiveness of thepresent invention at controlling the residual stresses.

The present invention allows the imparted strain levels to be accuratelycontrolled. The present invention is applicable to components with allmicrostructures commonly found in nickel base superalloy components,e.g. fine grains, medium grains, coarse grains or dual microstructures.The present invention is applicable to high strength nickel basesuperalloys for example RR1000, U720Li, Rene 95, Rene 88DT, ME3, N18,Alloy 10 and LSHR.

U720Li consists of 15 wt % cobalt, 16 wt % chromium, 3 wt % molybdenum,1.25 wt % tungsten, 5 wt % titanium, 2.5 wt % aluminium, 0.015 wt %boron, 0.015 wt % carbon and the balance nickel plus incidentalimpurities.

Rene 95 consists of 8.12 wt % cobalt, 12.94 wt % chromium, 3.45 wt %molybdenum, 3.43 wt % tungsten, 2.44 wt % titanium, 3.42 wt % aluminium,3.37 wt % niobium, 0.012 wt % boron, 0.05 wt % zirconium, 0.07 wt %carbon and the balance nickel plus incidental impurities.

Rene 88DT consists of 13.1 wt % cobalt, 15.8 wt % chromium, 4 wt %molybdenum, 3.9 wt % tungsten, 3.7 wt % titanium, 2 wt % aluminium, 0.7wt % niobium, 0.016 wt % boron, 0.045 wt % zirconium, 0.05 wt % carbonand the balance nickel plus incidental impurities.

ME3 consists of 20.6 wt % cobalt, 13 wt % chromium, 3.8 wt % molybdenum,2.1 wt % tungsten, 2.4 wt % tantalum, 3.7 wt % titanium, 3.4 wt %aluminium, 0.03 wt % boron, 0.05 wt % zirconium, 0.04 wt % carbon andthe balance nickel plus incidental impurities.

N18 consists of 15.4 wt % cobalt, 11.1 wt % chromium, 6.44 wt %molybdenum, 4.28 wt % titanium, 4.28 wt % aluminium, 0.5 wt % hafnium,0.008 wt % boron, 0.019 wt % zirconium, 0.022 wt % carbon and thebalance nickel plus incidental impurities.

Alloy 10 consists of 17.93 wt % cobalt, 10.46 wt % chromium, 2.52 wt %molybdenum, 4.74 wt % tungsten, 1.61 wt % tantalum, 3.79 wt % titanium,3.53 wt % aluminium, 0.028 wt % boron, 0.07 wt % zirconium, 0.027 wt %carbon and the balance nickel plus incidental impurities.

LSHR consists of 20.8 wt % cobalt, 12.7 wt % chromium, 2.74 wt %molybdenum, 4.37 wt % tungsten, 1.65 wt % tantalum, 3.47 wt % titanium,3.48 wt % aluminium, 0.028 wt % boron, 0.049 wt % zirconium, 0.024 wt %carbon and the balance nickel plus incidental impurities.

The present invention is also applicable to titanium base alloys, forexample Ti6246, Ti6242 or other alloys where increased tensileproperties or creep properties are required.

The present invention may be used to reduce, or eliminate, residualstresses developed by the solution heat treatment process. The presentinvention may be used to produce unique residual stress profiles in acomponent. The present invention may be used to support increasedprecipitation kinetics if it is applied before the ageing. The presentinvention may be used to selectively alter the retained strain orprecipitation kinetics within a superalloy disc. The present inventionincreases the mechanical strength of the alloy component by introducingdislocations, structural disturbances to the crystal structure, which inturn present obstacles to the creation and movement of furtherdislocations and hence increases mechanical strength.

The present invention enables turbine discs, compressor discs,compressor cones or turbine cover plates to be produced with enhancedproof and tensile strength and creep properties or reduced residualstress levels. This enables an increase in the operating life of thecomponent, enables an increase in the operating rotational speed of thecomponent, enables a decrease in the size of the component for anidentical gas turbine engine cycle or enables a reduction in weight ofthe component for the same operating life. The improved properties allowan increase in overspeed capability.

1. A method of improving the mechanical properties of a componentcomprising the steps of:— a) forging a preform to produce a shapedpreform with a predetermined shape at a first predetermined temperature,b) solution heat treating the shaped preform, c) quenching the shapedpreform, d) ageing the shaped preform, e) machining the shaped preformto a finished shape or a semi-finished shape, and f) forging the shapedpreform at a second predetermined temperature to impart a predeterminedresidual strain in the shaped preform after step c) and before step e),wherein the second predetermined temperature is less than the firstpredetermined temperature.
 2. A method as claimed in claim 1 whereinstep f) is after step c) and before step d).
 3. A method as claimed inclaim 1 wherein step f) is after step d) and before step e).
 4. A methodas claimed in claim 1 wherein step f) is concurrent with step d).
 5. Amethod as claimed in claim 1 wherein the second predeterminedtemperature is between 700° C. and 870° C.
 6. A method as claimed inclaim 1 wherein the second predetermined temperature is between 750° C.and 850° C.
 7. A method as claimed in claim 1 wherein the secondpredetermined temperature is between 760° C. and 810° C.
 8. A method asclaimed in claim 1 wherein the forging step f) imparts a predeterminedresidual tensile strain or a predetermined residual compressive strain.9. A method as claimed in claim 1 wherein the forging step f) imparts astrain of less than 10%.
 10. A method as claimed in claim 1 wherein themethod comprises machining the shaped preform after step a) and beforestep b).
 11. A method as claimed in claim 1 wherein step f) comprisesisothermally forging.
 12. A method as claimed in claim 1 wherein step f)comprises forging at a strain rate between 1×10⁻⁴ and 1×10⁻² s⁻¹.
 13. Amethod as claimed in claim 1 wherein step a) comprises isothermallyforging.
 14. A method as claimed in claim 1 wherein in step a) the firstpredetermined temperature is up to gamma prime solvus minus 25° C. to50° C.
 15. A method as claimed in claim 14 wherein step a) comprisesforging at a strain rate between 1×10⁻⁴ and 1×10⁻² s⁻¹.
 16. A method asclaimed in claim 1 wherein in step a) the first predeterminedtemperature is up to gamma prime solvus minus 55° C. to 110° C.
 17. Amethod as claimed in claim 16 wherein step a) comprises forging at astrain rate between 1×10⁻² and 5×10⁻¹ s⁻¹.
 18. A method as claimed inclaim 1 wherein step b) comprises a subsolvus solution heat treatmentand or a supersolvus heat treatment.
 19. A method as claimed in claim 18wherein step b) comprises a subsolvus solution heat treatment at 1120°C. for 4 hours.
 20. A method as claimed in claim 18 wherein step b)comprises a subsolvus solution heat treatment at 1120° C. for 4 hours,followed by quenching, followed by a supersolvus heat treatment at 1204°C. for 1 hour.
 21. A method as claimed in claim 18 wherein step b)comprises a supersolvus heat treatment at 1204° C. for 1 hour.
 22. Amethod as claimed in claim 1 wherein step d) comprises an ageing heattreatment at 760° C. for 16 hours.
 23. A method as claimed in claim 1wherein the component is selected from the group consisting of acompressor disc, a turbine disc, a compressor cone and a turbine coverplate.
 24. A method as claimed in claim 1 wherein the component isselected from the group consisting of a nickel base superalloy and atitanium base alloy.